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Low-temperature superplasticity and internal friction in microcrystalline Mg alloys processed by ECAP
Scripta Materialia 50 (2004) 861–865 www.actamat-journals.com
Low-temperature superplasticity and internal friction in microcrystalline Mg alloys processed by ECAP V.N. Chuvil’deev a, T.G. Nieh a
b,*
, M.Yu. Gryaznov a, A.N. Sysoev a, V.I. Kopylov
c
Research Physical & Technical Institute, Nizhny Novgorod State University, 23/3 Gagarin Ave., Nizhny Novgorod 603950, Russia b Lawrence Livermore National Laboratory, P.O. Box 808, L-350, Livermore, CA 94551-9900, USA c Physical–Technical Institute of National Academy of Science, Kuprievicha Str. 10, Minsk 220141, Belarus Received 14 August 2003; received in revised form 27 October 2003; accepted 8 December 2003
Abstract Excellent low-temperature superplasticity (<300 C) was observed in the ECAP-processed materials: elongations to failure are 810% and 570% at a strain rate of 3 · 103 s1 for ZK60 and AZ91 alloys, respectively. Internal friction can be practically used to determine the optimum temperature for superplasticity. 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mg alloys; Severe plastic deformation; Low-temperature superplasticity; Internal friction
1. Introduction Mg alloys are one of the lightest structural materials. However, they have relatively low plasticity, which prevents a wide application of Mg alloys. For commercial cast Mg alloys an elongation to failure is usually less than 20% at room temperature and is about 30–40% at 300 C [1]. A practical way to enhance ductility of alloys is to create a microstructure that would facilitate the occurrence of superplasticity. Thus, low-temperature superplasticity (LTSP) of Mg alloys (from room temperature to 300 C) is of special interest. LTSP has been demonstrated in many alloys with non-equilibrium microcrystalline structures processed by equal-channel angular pressing (ECAP) [2–4]. Particularly, LTSP has been reported in some microcrystalline Mg alloys processed by ECAP at relatively low strain rates (104 –105 s1 ) [5–12]. For commercial interest, however, it is desirable to obtain superplasticity in high-strength Mg alloys not only at low temperatures but also at high strain rates ( > 3 · 103 s1 ). To achieve this goal in highstrength ZK60 and AZ91 Mg alloys, two major tasks are encountered. The first task is to optimize the ECAP process to develop ultrafine grains with non-equilibrium grain boundaries [2,4]. The second task is to optimize *
Corresponding author. E-mail address: [email protected] (T.G. Nieh).
the superplasticity properties. In the present study, we demonstrate that internal friction method can be effectively used to identify the optimum temperature range for superplasticity in ZK60 and AZ91 Mg alloys.
2. Experiments Two Mg alloys: ZK60 (Mg–6wt.%Zn–0.5wt.%Zr) and AZ91 (Mg–9wt.%Al–1wt.%Zn–0.2wt.%Mn), both in cast and microcrystalline states, were studiedin the present work. Cast alloys have microstructures with inhomogeneous grain size distribution, and the mean grain sizes in cast ZK60 and AZ91 alloys were 20 and 15 lm, respectively. Microcrystalline alloys were obtained by ECAP processing of the cast samples with an initial size of 22 mm · 22 mm · 160 mm. Multiple pressings through a channel die with a 90 die angle were performed 6 times with a velocity of 0.4 mm/s at temperatures between 180 and 230 C for ZK60, and between 200 and 250 C for AZ91 using the so-called ‘‘route Bc’’ process [13]. As a result of ECAP, homogeneous microcrystalline structures with a respective mean grain size of 1 and 0.8 lm in ZK60 and AZ91, were produced. High strain rate superplasticity tests were performed using an automated 2167 P-50 (analogous to Instron machine). Tensile samples were mechanically polished
1359-6462/$ - see front matter 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2003.12.003
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before testing and then spark eroded to produce a square cross-section of 2 mm · 2 mm with a gauge length of 3 mm. Superplasticity tests were carried out at 200– 430 C in an air furnace and initial strain rate ranging from 6 · 104 to 3 · 101 s1 . To attain thermal equilibrium, samples were held at the testing temperature for 400 s before pulling. An inverted torsion pendulum was used to measure internal friction. It allows the study of grain boundary internal friction under the conditions of free damped oscillations in the frequency range of 1–10 Hz. Samples with a dimension of 1.5 mm · 1.5 mm · 50 mm were used for measurements conducted at a frequency of 5 Hz. The strain amplitude was limited to 5 · 106 . Continuous heating and cooling of the test samples were carried out in the temperature range of 20–350 C with a constant rate of 4 K/min.
3. Results 3.1. Internal friction
Internal friction, * 10
-3
Temperature dependence of internal friction Q1 ðT Þ at a frequency of 5 Hz during heating and cooling of microcrystalline ZK60 and AZ91 alloys is presented in Fig. 1. During heating from 20 to 260 C, internal friction in the microcrystalline ZK60 alloy increases monotonically from 3.5 · 103 to a high value of about 3.2 · 101 . Further heating to 260–325 C causes internal friction to decrease sharply to a value of 5 · 102 . A similar but less pronounced response is observed in the microcrystalline AZ91 alloy. In this case, internal friction increases rapidly from 2.5 · 103 to 1.6 · 101 when the sample is heated from 20 to 300 C (Fig. 1). The internal-friction peak is also less pronounced and shifted 40 C toward high temperatures, as compared to that in the microcrystalline ZK60 alloy. For comparison, data for cast AZ91 are also included in Fig. 1. For the cast AZ91, no internal-friction peak is observed during either
ECAP ZK60 (heating) ECAP AZ91 (heating) ECAP AZ91 (cooling) Cast AZ91 (heating)
300
heating and cooling. It is evident in the Fig. 1 that the cooling curves of Q1 ðT Þ for both microcrystalline alloys are similar to that for the cast samples. It is noted that the internal friction peaks observed during the first heating is absent during cooling and consequent reheating. This indicates the irreversible nature of these peaks. 3.2. Low-temperature superplasticity Tensile stress–strain curves for microcrystalline ZK60 and AZ91 at temperatures between 200 and 430 C and a fixed strain rate of 3 · 103 s1 are shown in Fig. 2a and b, respectively. An increase in ductility and decrease in flow stress is observed between 200 and 260 C. The temperature of 260 C is apparently the optimum temperature for microcrystalline ZK60 as its stress–strain curve exhibits the lowest flow stress and maximum ductility. A higher flow stress and lower ductility observed at 280 C is probably resulted from the fact that excessive grain growth takes place in the microcrystalline ZK60 at T > 260 C. For microcrystalline AZ91 alloy this occurs at a slightly higher temperature of 300 C. Fig. 3 displays the elongation to failure, d, as a function of deformation temperature for both microcrystalline ZK60 and AZ91. As shown in the figure, tensile elongation as a function of test temperature at the fixed strain rate of 3 · 103 s1 is non-monotonic. In the microcrystalline ZK60 alloy, the elongation increases rapidly from 170% to 810% as the deformation temperature increases from 150 to 260 C. However, it decreases sharply down to 170% at 280 C (curve 1 in Fig. 3). A further increase to 430 C causes a slightly improved plasticity to 340%. A similar behavior is observed in the microcrystalline AZ91 alloy where d increases monotonically from 30% to 570% as the temperature increases from 200 to 300 C (curve 2 in Fig. 3). Above 325 C, however, d decreases rapidly to 350%. The elongation to failure and ultimate tensile strength as functions of strain rate, e_ , at the optimal superplasticity temperature (260 C for ZK60 and 300 C for AZ91) are presented in Fig. 4. With a decreasing strain rate, the values of d increase and reach 960% and 570% respectively in ZK60 and AZ91 alloys at e_ ¼ 6:5 104 s1 and e_ ¼ 3 103 s1 .
200 100
4. Analysis
0 0
100
200
300
400
o
Temperature, C Fig. 1. Temperature dependence of internal friction in microcrystalline ZK60 and AZ91, and cast AZ91 Mg alloys (frequency of 5 Hz). The cooling curves for ECAP ZK60 and AZ91 essentially overlap.
A comparison between the data obtained in the present work and those from previous studies [1,6–8] is summarized in Table 1. The data show that the ductility of microcrystalline ZK60 and AZ91 alloys at temperatures range of 200–300 C and at a strain rate of about 103 s1 is almost twice higher than that for similar al-
30
280°c 200°c
20
240°c 10
430°c
260°c
863
60
ECAP ZK60, 0.003 1/s
40
Nominal stress, MPa
Nominal stress, MPa
V.N. Chuvil’deev et al. / Scripta Materialia 50 (2004) 861–865
0
ECAP AZ91, 0.003 1/s 50
200°c
40 250°c
30
300°c
20 350°c
10
400°c
0 0
100
(a)
200
300
400
500
600
700
0
800
100
(b)
Nominal strain, %
200
300
400
500
600
Nominal strain, %
Fig. 2. Nominal stress–strain relations at strain rate of 3 · 103 s1 for microcrystalline ZK60 (a) and AZ91 (b) Mg alloys.
Elongation to failure, %
1000
room temperature and a strain rate 3 · 103 s1 (compared to merely 5–10% in cast AZ91 [10]). The improved room temperature ductility was suggested to be a result of texture randomization [12]. It is especially noted that with the ductility improvement the strength of ECAPprocessed materials was not sacrificed. From an application point of view, the effect of ductility enhancement at room temperature without a significant decrease in strength (Table 1) is obviously very desirable.
ECAP ZK60 (0.003 1/s) ECAP AZ91 (0.003 1/s)
800
1
600
2
400 200 0 0
100
200
300
400
5. Discussion
Temperature, oc
1200
ECAP ZK60 (elongation) ECAP ZK60 (UTS)
84
ECAP AZ91 (elongation) ECAP AZ91 (UTS)
1000
70
800
56
600
42
400
28
200
14
0 0.0001
UTS, MPa
Elongation to failure, %
Fig. 3. Elongation to failure as a function of deformation temperature for microcrystalline ZK60 (1) and AZ91 (2) alloys at a strain rate of 3 · 103 s1 .
0 0.001
0.01
0.1
1
Strain rate, s-1
Fig. 4. Elongation to failure and ultimate tensile strength (UTS) as functions of strain rate for microcrystalline ZK60 (260 C) and AZ91 (300 C).
loys not produced by ECAP. It is also evident in Table 1 that ECAP not only enhances high temperature plasticity in ZK60 and AZ91 but also improves the room temperature properties of the alloys. For instance, elongation to failure for the microcrystalline ZK60 alloy at room temperature averages about 45%, which is more than twice the value for the same alloy processed differently [1]. Also, the microcrystalline AZ91 alloy exhibits a significantly improved 30% elongation at
To illustrate the practicality of internal friction for superplasticity study, the temperature dependences of both internal friction and elongation to failure as a function of deformation temperature in microcrystalline ZK60 and AZ91 alloys are shown in Fig. 5a and b. It is evident from the figure that the two curves not only show a similar trend but also the peaks of the two curves occur at essentially the same temperature. Since the internal friction peak results from grain boundary sliding, the fact that the peak positions for both internal friction and elongation to failure coincide with each other (Fig. 5a and b) suggests that grain boundary sliding is the dominant deformation mechanism for superplasticity in both microcrystalline ZK60 and AZ91 alloys. The constitutive law for superplastic flow is expressed as [14] r n b p GD u Q b0 e_ ¼ A exp ; ð1Þ G d kT RT where b is Burgers vector, d is mean grain size, u is thickness of grain boundary (which is about 0.5–1 nm), Db0 is pre-exponential factor of grain-boundary diffusion, G is shear modulus, k is Boltzmann constant, Q is activation energy of grain-boundary diffusion, R is molar gas constant, p ¼ 2 and A ¼ 102 are the numerical factors. The parameter n is the reciprocal of strain-rate sensitivity coefficient ðm ¼ @ ln r=@ ln e_ Þ. The n values can be determined from the analysis of the experimental
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Table 1 Tensile properties at room temperature and temperatures between 150 and 325 C for ZK60 and AZ91 alloys processed using different techniques Deformation temperature (C)
Elongation-tofailure (%)
Ultimate tensile strength (MPa)
Ref.
ZK60 ZK60 ZK60 ZK60 ZK60 ZK60 ZK60 AZ91 AZ91 AZ91 AZ91 AZ91 AZ91 AZ91
ECAP ECAP Extrusion Forging ECAP Extrusion Casting ECAP ECAP ECAP Extrusion ECAP Casting Extrusion
3 · 103 1 · 103 1 · 103 – 3 · 103 – – 3 · 102 3 · 103 6 · 105 6 · 105 3 · 103 – –
260 200 325 150 20 20 20 250 300 200 200 20 20 20
810 420 400 36 45 18 5 375 570 660 80 30 3 13
12 30 16 190 260 370 275 55 23 25 100 250 235 341
This [7] [6] [1] This [1] [1] This This [8] [8] This [10] [10]
600
220
300
110
0 0
(a)
100
200
300
ECAP AZ91 400
140
200
70
0 0
100
(b)
o
Temperature, C
work work
work
210
600
0 400
work
-3
330 ECAP ZK60
work
200
300
Internal friction, *10
900
Elongation to failure,%
Strain rate (1/s)
-3 Internal friction, *10
Production method
Elongation to failure,%
Alloy
0 400
o
Temperature, C
Fig. 5. Comparison of elongation to failure as a function of deformation temperature and temperature dependence of internal friction in microcrystalline ZK60 (a) and AZ91 (b) alloys.
dependence rð_eÞ, and they are n ¼ 2:5 ðm ¼ 0:4Þ and n ¼ 3:3 ðm ¼ 0:3Þ for microcrystalline ZK60 and AZ91, respectively. From Eq. (1) the activation energies can be determined to be 75 and 78 kJ/mol for microcrystalline ZK60 and AZ91 alloys, respectively. These values are much lower than the reported value of 92 kJ/mol for the grain-boundary diffusion in pure Mg [15] and Mg alloys [6,16,17] with equilibrium grain boundaries. As a result of low activation energy, the grain-boundary diffusion coefficient Db ¼ Db0 expðQ=RT Þ at 260 C is 25–40 times higher than the value of Db ¼ 7:4 1012 m2 /s calculated on the basis of Q ¼ 92 kJ/mol. This accelerated diffusion explains the occurrence of superplasticity in microcrystalline Mg alloys at relatively low temperatures of 200–300 C and strain rate of 3 · 103 s1 . A model to explain the apparently enhanced grain boundary diffusion in a material containing non-equilibrium grain boundary structures during superplastic deformation has been proposed [4,18]. In accordance with the model [18], a value of Db ð_eÞ can be written as: Db ð_eÞ ¼ Db ð1 þ M1 lnð1 þ M2 e_ ÞÞ;
ð2Þ
where parameters M1 and M2 can be calculated from the models [4] and are M1 ¼ 150 and M2 ¼ 90 for Mg at T ¼ 260 C. Substituting these values into expression for Db ð_eÞ, one can readily obtain the ratio Db ð_eÞ=Db , grainboundary diffusion enhancement factor, to be equal to 35 at e_ ¼ 3 103 s1 . This enhancement is sufficient to accommodate strain at grain triple junctions during grain boundary sliding. Otherwise, cavitation would immediately occur and cause premature fracture. Therefore, by taking into consideration of non-equilibrium state of grain boundaries, the achievement of superplasticity in microcrystalline ZK60 and AZ91 at low temperatures of 200–300 C can be reasonably explained.
6. Conclusions In the present study, we demonstrated the following: 1. ECAP can produce a homogeneous microcrystalline structure with non-equilibrium grain boundaries in
V.N. Chuvil’deev et al. / Scripta Materialia 50 (2004) 861–865
ZK60 and AZ91 alloys. As a result, low-temperature superplasticity at strain rate of 3 · 103 s1 : elongation to failure of 810% (260 C) and 570% (300 C) respectively in ZK60 and AZ91 alloys, can be achieved. 2. It is shown that the temperature at which internalfriction peak occurs coincides with the temperature for maximum tensile elongation, suggesting internal friction is a practical method for determining the optimum temperature for superplasticity. 3. The activation energies for superplasticity are 75 and 78 kJ/mol for the ECAP-processed ZK60 and AZ91 alloys, respectively. These values are lower than 92 kJ/mol reported in the literature for the grain boundary diffusion in Mg. The enhanced grain boundary diffusion is caused by the fact that ECAP-processed alloys possess non-equilibrium grain boundaries. This grain boundary diffusion enhancement greatly assists strain accommodation at grain triple junctions during superplastic deformation and promotes low-temperature superplasticity.
Acknowledgements The authors (VNC, MYG and ANS) thank the Russian Foundation of Basic Research (Grants: 02-0333043, 03-02-16923), Ministry of Education of Russian Federation (Grant: E02-4.0-131) and the Basic Research and Higher Education (BRHE) program and Research and Education Center of ‘‘Physics of solid-state nanostructures’’ at Nizhny Novgorod State University for financial and technical support. This work was also performed under the auspices of the US Department of Energy by the University of California, Lawrence Liv-
865
ermore National Laboratory under Contract no. W7405-Eng-48. References [1] Avedisian M, Baker H, editors. ASM specialty handbook mg and mg alloys. Materials Park, Ohio: ASM International; 1999. [2] Segal VM, Kopylov VI, Reznikov VI, Pavlik DA, Malyshev VF. Processes of plastic structure formation of metals. Minsk: Nauka i Tehnika; 1994. [3] Kopylov VI. In: Investigations and applications of severe plastic deformation.Lowe TC, Valiev RZ, editors. NATO ASI series 3 High technology. Netherlands: Kluwer Academic Publisher; 2000. p. 23. [4] Chuvil’deev VN, Kopylov VI, Zeiger W. Annales de Chim Sci des Mater 2002;27:55. [5] Chuvil’deev VN, Kopylov VI, Gryaznov MY, Sysoev AN. Doklady Phys 2003;48:343. [6] Watanabe H, Mukai T, Higashi K. Scripta Mater 1999;40:477. [7] Watanabe H, Mukai T, Ishikawa K, Higashi K. Scripta Mater 2002;46:851. [8] Mabuchi M, Iwasaki H, Yanase K, Higashi K. Scripta Mater 1997;36:681. [9] Mabuchi M, Ameyama K, Iwasaki H, Higashi K. Acta Mater 1999;47:2047. [10] Kubota K, Mabuchi M, Higashi K. J Mater Sci 1999;34:2255. [11] Kamado S, Ashie T, Yamada H. Mater Sci Forum 2000;350– 351:65. [12] Mukai T, Yamanoi M, Watanabe H, Higashi K. Scripta Mater 2001;45:89. [13] Furukawa M, Iwahashi Y, Horita Z, Nemoto M, Langdon TG. Mater Sci Eng A 1998;257:328. [14] Nieh TG, Wadsworth J, Sherby OD. Superplasticity in metals and ceramics. Cambridge: Cambridge University Press; 1997. [15] Frost HJ, Ashby MF. Deformation-mechanisms maps. Oxford: Pergamon Press; 1982. p. 166. [16] Galiyev A, Kaibyshev R, Gottstein G. Acta Mater 2001;49:1199. [17] Watanabe H, Mukai T, Kohzu M, Tanabe S, Higashi K. Acta Mater 1999;47:3753. [18] Chuvil’deev VN, Petryaev AV. Phys Metals Metallogr 2000;89:122.
Low-temperature superplasticity and internal friction in microcrystalline Mg alloys processed by ECAP V.N. Chuvil’deev a, T.G. Nieh a
b,*
, M.Yu. Gryaznov a, A.N. Sysoev a, V.I. Kopylov
c
Research Physical & Technical Institute, Nizhny Novgorod State University, 23/3 Gagarin Ave., Nizhny Novgorod 603950, Russia b Lawrence Livermore National Laboratory, P.O. Box 808, L-350, Livermore, CA 94551-9900, USA c Physical–Technical Institute of National Academy of Science, Kuprievicha Str. 10, Minsk 220141, Belarus Received 14 August 2003; received in revised form 27 October 2003; accepted 8 December 2003
Abstract Excellent low-temperature superplasticity (<300 C) was observed in the ECAP-processed materials: elongations to failure are 810% and 570% at a strain rate of 3 · 103 s1 for ZK60 and AZ91 alloys, respectively. Internal friction can be practically used to determine the optimum temperature for superplasticity. 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Mg alloys; Severe plastic deformation; Low-temperature superplasticity; Internal friction
1. Introduction Mg alloys are one of the lightest structural materials. However, they have relatively low plasticity, which prevents a wide application of Mg alloys. For commercial cast Mg alloys an elongation to failure is usually less than 20% at room temperature and is about 30–40% at 300 C [1]. A practical way to enhance ductility of alloys is to create a microstructure that would facilitate the occurrence of superplasticity. Thus, low-temperature superplasticity (LTSP) of Mg alloys (from room temperature to 300 C) is of special interest. LTSP has been demonstrated in many alloys with non-equilibrium microcrystalline structures processed by equal-channel angular pressing (ECAP) [2–4]. Particularly, LTSP has been reported in some microcrystalline Mg alloys processed by ECAP at relatively low strain rates (104 –105 s1 ) [5–12]. For commercial interest, however, it is desirable to obtain superplasticity in high-strength Mg alloys not only at low temperatures but also at high strain rates ( > 3 · 103 s1 ). To achieve this goal in highstrength ZK60 and AZ91 Mg alloys, two major tasks are encountered. The first task is to optimize the ECAP process to develop ultrafine grains with non-equilibrium grain boundaries [2,4]. The second task is to optimize *
Corresponding author. E-mail address: [email protected] (T.G. Nieh).
the superplasticity properties. In the present study, we demonstrate that internal friction method can be effectively used to identify the optimum temperature range for superplasticity in ZK60 and AZ91 Mg alloys.
2. Experiments Two Mg alloys: ZK60 (Mg–6wt.%Zn–0.5wt.%Zr) and AZ91 (Mg–9wt.%Al–1wt.%Zn–0.2wt.%Mn), both in cast and microcrystalline states, were studiedin the present work. Cast alloys have microstructures with inhomogeneous grain size distribution, and the mean grain sizes in cast ZK60 and AZ91 alloys were 20 and 15 lm, respectively. Microcrystalline alloys were obtained by ECAP processing of the cast samples with an initial size of 22 mm · 22 mm · 160 mm. Multiple pressings through a channel die with a 90 die angle were performed 6 times with a velocity of 0.4 mm/s at temperatures between 180 and 230 C for ZK60, and between 200 and 250 C for AZ91 using the so-called ‘‘route Bc’’ process [13]. As a result of ECAP, homogeneous microcrystalline structures with a respective mean grain size of 1 and 0.8 lm in ZK60 and AZ91, were produced. High strain rate superplasticity tests were performed using an automated 2167 P-50 (analogous to Instron machine). Tensile samples were mechanically polished
1359-6462/$ - see front matter 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.scriptamat.2003.12.003
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V.N. Chuvil’deev et al. / Scripta Materialia 50 (2004) 861–865
before testing and then spark eroded to produce a square cross-section of 2 mm · 2 mm with a gauge length of 3 mm. Superplasticity tests were carried out at 200– 430 C in an air furnace and initial strain rate ranging from 6 · 104 to 3 · 101 s1 . To attain thermal equilibrium, samples were held at the testing temperature for 400 s before pulling. An inverted torsion pendulum was used to measure internal friction. It allows the study of grain boundary internal friction under the conditions of free damped oscillations in the frequency range of 1–10 Hz. Samples with a dimension of 1.5 mm · 1.5 mm · 50 mm were used for measurements conducted at a frequency of 5 Hz. The strain amplitude was limited to 5 · 106 . Continuous heating and cooling of the test samples were carried out in the temperature range of 20–350 C with a constant rate of 4 K/min.
3. Results 3.1. Internal friction
Internal friction, * 10
-3
Temperature dependence of internal friction Q1 ðT Þ at a frequency of 5 Hz during heating and cooling of microcrystalline ZK60 and AZ91 alloys is presented in Fig. 1. During heating from 20 to 260 C, internal friction in the microcrystalline ZK60 alloy increases monotonically from 3.5 · 103 to a high value of about 3.2 · 101 . Further heating to 260–325 C causes internal friction to decrease sharply to a value of 5 · 102 . A similar but less pronounced response is observed in the microcrystalline AZ91 alloy. In this case, internal friction increases rapidly from 2.5 · 103 to 1.6 · 101 when the sample is heated from 20 to 300 C (Fig. 1). The internal-friction peak is also less pronounced and shifted 40 C toward high temperatures, as compared to that in the microcrystalline ZK60 alloy. For comparison, data for cast AZ91 are also included in Fig. 1. For the cast AZ91, no internal-friction peak is observed during either
ECAP ZK60 (heating) ECAP AZ91 (heating) ECAP AZ91 (cooling) Cast AZ91 (heating)
300
heating and cooling. It is evident in the Fig. 1 that the cooling curves of Q1 ðT Þ for both microcrystalline alloys are similar to that for the cast samples. It is noted that the internal friction peaks observed during the first heating is absent during cooling and consequent reheating. This indicates the irreversible nature of these peaks. 3.2. Low-temperature superplasticity Tensile stress–strain curves for microcrystalline ZK60 and AZ91 at temperatures between 200 and 430 C and a fixed strain rate of 3 · 103 s1 are shown in Fig. 2a and b, respectively. An increase in ductility and decrease in flow stress is observed between 200 and 260 C. The temperature of 260 C is apparently the optimum temperature for microcrystalline ZK60 as its stress–strain curve exhibits the lowest flow stress and maximum ductility. A higher flow stress and lower ductility observed at 280 C is probably resulted from the fact that excessive grain growth takes place in the microcrystalline ZK60 at T > 260 C. For microcrystalline AZ91 alloy this occurs at a slightly higher temperature of 300 C. Fig. 3 displays the elongation to failure, d, as a function of deformation temperature for both microcrystalline ZK60 and AZ91. As shown in the figure, tensile elongation as a function of test temperature at the fixed strain rate of 3 · 103 s1 is non-monotonic. In the microcrystalline ZK60 alloy, the elongation increases rapidly from 170% to 810% as the deformation temperature increases from 150 to 260 C. However, it decreases sharply down to 170% at 280 C (curve 1 in Fig. 3). A further increase to 430 C causes a slightly improved plasticity to 340%. A similar behavior is observed in the microcrystalline AZ91 alloy where d increases monotonically from 30% to 570% as the temperature increases from 200 to 300 C (curve 2 in Fig. 3). Above 325 C, however, d decreases rapidly to 350%. The elongation to failure and ultimate tensile strength as functions of strain rate, e_ , at the optimal superplasticity temperature (260 C for ZK60 and 300 C for AZ91) are presented in Fig. 4. With a decreasing strain rate, the values of d increase and reach 960% and 570% respectively in ZK60 and AZ91 alloys at e_ ¼ 6:5 104 s1 and e_ ¼ 3 103 s1 .
200 100
4. Analysis
0 0
100
200
300
400
o
Temperature, C Fig. 1. Temperature dependence of internal friction in microcrystalline ZK60 and AZ91, and cast AZ91 Mg alloys (frequency of 5 Hz). The cooling curves for ECAP ZK60 and AZ91 essentially overlap.
A comparison between the data obtained in the present work and those from previous studies [1,6–8] is summarized in Table 1. The data show that the ductility of microcrystalline ZK60 and AZ91 alloys at temperatures range of 200–300 C and at a strain rate of about 103 s1 is almost twice higher than that for similar al-
30
280°c 200°c
20
240°c 10
430°c
260°c
863
60
ECAP ZK60, 0.003 1/s
40
Nominal stress, MPa
Nominal stress, MPa
V.N. Chuvil’deev et al. / Scripta Materialia 50 (2004) 861–865
0
ECAP AZ91, 0.003 1/s 50
200°c
40 250°c
30
300°c
20 350°c
10
400°c
0 0
100
(a)
200
300
400
500
600
700
0
800
100
(b)
Nominal strain, %
200
300
400
500
600
Nominal strain, %
Fig. 2. Nominal stress–strain relations at strain rate of 3 · 103 s1 for microcrystalline ZK60 (a) and AZ91 (b) Mg alloys.
Elongation to failure, %
1000
room temperature and a strain rate 3 · 103 s1 (compared to merely 5–10% in cast AZ91 [10]). The improved room temperature ductility was suggested to be a result of texture randomization [12]. It is especially noted that with the ductility improvement the strength of ECAPprocessed materials was not sacrificed. From an application point of view, the effect of ductility enhancement at room temperature without a significant decrease in strength (Table 1) is obviously very desirable.
ECAP ZK60 (0.003 1/s) ECAP AZ91 (0.003 1/s)
800
1
600
2
400 200 0 0
100
200
300
400
5. Discussion
Temperature, oc
1200
ECAP ZK60 (elongation) ECAP ZK60 (UTS)
84
ECAP AZ91 (elongation) ECAP AZ91 (UTS)
1000
70
800
56
600
42
400
28
200
14
0 0.0001
UTS, MPa
Elongation to failure, %
Fig. 3. Elongation to failure as a function of deformation temperature for microcrystalline ZK60 (1) and AZ91 (2) alloys at a strain rate of 3 · 103 s1 .
0 0.001
0.01
0.1
1
Strain rate, s-1
Fig. 4. Elongation to failure and ultimate tensile strength (UTS) as functions of strain rate for microcrystalline ZK60 (260 C) and AZ91 (300 C).
loys not produced by ECAP. It is also evident in Table 1 that ECAP not only enhances high temperature plasticity in ZK60 and AZ91 but also improves the room temperature properties of the alloys. For instance, elongation to failure for the microcrystalline ZK60 alloy at room temperature averages about 45%, which is more than twice the value for the same alloy processed differently [1]. Also, the microcrystalline AZ91 alloy exhibits a significantly improved 30% elongation at
To illustrate the practicality of internal friction for superplasticity study, the temperature dependences of both internal friction and elongation to failure as a function of deformation temperature in microcrystalline ZK60 and AZ91 alloys are shown in Fig. 5a and b. It is evident from the figure that the two curves not only show a similar trend but also the peaks of the two curves occur at essentially the same temperature. Since the internal friction peak results from grain boundary sliding, the fact that the peak positions for both internal friction and elongation to failure coincide with each other (Fig. 5a and b) suggests that grain boundary sliding is the dominant deformation mechanism for superplasticity in both microcrystalline ZK60 and AZ91 alloys. The constitutive law for superplastic flow is expressed as [14] r n b p GD u Q b0 e_ ¼ A exp ; ð1Þ G d kT RT where b is Burgers vector, d is mean grain size, u is thickness of grain boundary (which is about 0.5–1 nm), Db0 is pre-exponential factor of grain-boundary diffusion, G is shear modulus, k is Boltzmann constant, Q is activation energy of grain-boundary diffusion, R is molar gas constant, p ¼ 2 and A ¼ 102 are the numerical factors. The parameter n is the reciprocal of strain-rate sensitivity coefficient ðm ¼ @ ln r=@ ln e_ Þ. The n values can be determined from the analysis of the experimental
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Table 1 Tensile properties at room temperature and temperatures between 150 and 325 C for ZK60 and AZ91 alloys processed using different techniques Deformation temperature (C)
Elongation-tofailure (%)
Ultimate tensile strength (MPa)
Ref.
ZK60 ZK60 ZK60 ZK60 ZK60 ZK60 ZK60 AZ91 AZ91 AZ91 AZ91 AZ91 AZ91 AZ91
ECAP ECAP Extrusion Forging ECAP Extrusion Casting ECAP ECAP ECAP Extrusion ECAP Casting Extrusion
3 · 103 1 · 103 1 · 103 – 3 · 103 – – 3 · 102 3 · 103 6 · 105 6 · 105 3 · 103 – –
260 200 325 150 20 20 20 250 300 200 200 20 20 20
810 420 400 36 45 18 5 375 570 660 80 30 3 13
12 30 16 190 260 370 275 55 23 25 100 250 235 341
This [7] [6] [1] This [1] [1] This This [8] [8] This [10] [10]
600
220
300
110
0 0
(a)
100
200
300
ECAP AZ91 400
140
200
70
0 0
100
(b)
o
Temperature, C
work work
work
210
600
0 400
work
-3
330 ECAP ZK60
work
200
300
Internal friction, *10
900
Elongation to failure,%
Strain rate (1/s)
-3 Internal friction, *10
Production method
Elongation to failure,%
Alloy
0 400
o
Temperature, C
Fig. 5. Comparison of elongation to failure as a function of deformation temperature and temperature dependence of internal friction in microcrystalline ZK60 (a) and AZ91 (b) alloys.
dependence rð_eÞ, and they are n ¼ 2:5 ðm ¼ 0:4Þ and n ¼ 3:3 ðm ¼ 0:3Þ for microcrystalline ZK60 and AZ91, respectively. From Eq. (1) the activation energies can be determined to be 75 and 78 kJ/mol for microcrystalline ZK60 and AZ91 alloys, respectively. These values are much lower than the reported value of 92 kJ/mol for the grain-boundary diffusion in pure Mg [15] and Mg alloys [6,16,17] with equilibrium grain boundaries. As a result of low activation energy, the grain-boundary diffusion coefficient Db ¼ Db0 expðQ=RT Þ at 260 C is 25–40 times higher than the value of Db ¼ 7:4 1012 m2 /s calculated on the basis of Q ¼ 92 kJ/mol. This accelerated diffusion explains the occurrence of superplasticity in microcrystalline Mg alloys at relatively low temperatures of 200–300 C and strain rate of 3 · 103 s1 . A model to explain the apparently enhanced grain boundary diffusion in a material containing non-equilibrium grain boundary structures during superplastic deformation has been proposed [4,18]. In accordance with the model [18], a value of Db ð_eÞ can be written as: Db ð_eÞ ¼ Db ð1 þ M1 lnð1 þ M2 e_ ÞÞ;
ð2Þ
where parameters M1 and M2 can be calculated from the models [4] and are M1 ¼ 150 and M2 ¼ 90 for Mg at T ¼ 260 C. Substituting these values into expression for Db ð_eÞ, one can readily obtain the ratio Db ð_eÞ=Db , grainboundary diffusion enhancement factor, to be equal to 35 at e_ ¼ 3 103 s1 . This enhancement is sufficient to accommodate strain at grain triple junctions during grain boundary sliding. Otherwise, cavitation would immediately occur and cause premature fracture. Therefore, by taking into consideration of non-equilibrium state of grain boundaries, the achievement of superplasticity in microcrystalline ZK60 and AZ91 at low temperatures of 200–300 C can be reasonably explained.
6. Conclusions In the present study, we demonstrated the following: 1. ECAP can produce a homogeneous microcrystalline structure with non-equilibrium grain boundaries in
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ZK60 and AZ91 alloys. As a result, low-temperature superplasticity at strain rate of 3 · 103 s1 : elongation to failure of 810% (260 C) and 570% (300 C) respectively in ZK60 and AZ91 alloys, can be achieved. 2. It is shown that the temperature at which internalfriction peak occurs coincides with the temperature for maximum tensile elongation, suggesting internal friction is a practical method for determining the optimum temperature for superplasticity. 3. The activation energies for superplasticity are 75 and 78 kJ/mol for the ECAP-processed ZK60 and AZ91 alloys, respectively. These values are lower than 92 kJ/mol reported in the literature for the grain boundary diffusion in Mg. The enhanced grain boundary diffusion is caused by the fact that ECAP-processed alloys possess non-equilibrium grain boundaries. This grain boundary diffusion enhancement greatly assists strain accommodation at grain triple junctions during superplastic deformation and promotes low-temperature superplasticity.
Acknowledgements The authors (VNC, MYG and ANS) thank the Russian Foundation of Basic Research (Grants: 02-0333043, 03-02-16923), Ministry of Education of Russian Federation (Grant: E02-4.0-131) and the Basic Research and Higher Education (BRHE) program and Research and Education Center of ‘‘Physics of solid-state nanostructures’’ at Nizhny Novgorod State University for financial and technical support. This work was also performed under the auspices of the US Department of Energy by the University of California, Lawrence Liv-
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ermore National Laboratory under Contract no. W7405-Eng-48. References [1] Avedisian M, Baker H, editors. ASM specialty handbook mg and mg alloys. Materials Park, Ohio: ASM International; 1999. [2] Segal VM, Kopylov VI, Reznikov VI, Pavlik DA, Malyshev VF. Processes of plastic structure formation of metals. Minsk: Nauka i Tehnika; 1994. [3] Kopylov VI. In: Investigations and applications of severe plastic deformation.Lowe TC, Valiev RZ, editors. NATO ASI series 3 High technology. Netherlands: Kluwer Academic Publisher; 2000. p. 23. [4] Chuvil’deev VN, Kopylov VI, Zeiger W. Annales de Chim Sci des Mater 2002;27:55. [5] Chuvil’deev VN, Kopylov VI, Gryaznov MY, Sysoev AN. Doklady Phys 2003;48:343. [6] Watanabe H, Mukai T, Higashi K. Scripta Mater 1999;40:477. [7] Watanabe H, Mukai T, Ishikawa K, Higashi K. Scripta Mater 2002;46:851. [8] Mabuchi M, Iwasaki H, Yanase K, Higashi K. Scripta Mater 1997;36:681. [9] Mabuchi M, Ameyama K, Iwasaki H, Higashi K. Acta Mater 1999;47:2047. [10] Kubota K, Mabuchi M, Higashi K. J Mater Sci 1999;34:2255. [11] Kamado S, Ashie T, Yamada H. Mater Sci Forum 2000;350– 351:65. [12] Mukai T, Yamanoi M, Watanabe H, Higashi K. Scripta Mater 2001;45:89. [13] Furukawa M, Iwahashi Y, Horita Z, Nemoto M, Langdon TG. Mater Sci Eng A 1998;257:328. [14] Nieh TG, Wadsworth J, Sherby OD. Superplasticity in metals and ceramics. Cambridge: Cambridge University Press; 1997. [15] Frost HJ, Ashby MF. Deformation-mechanisms maps. Oxford: Pergamon Press; 1982. p. 166. [16] Galiyev A, Kaibyshev R, Gottstein G. Acta Mater 2001;49:1199. [17] Watanabe H, Mukai T, Kohzu M, Tanabe S, Higashi K. Acta Mater 1999;47:3753. [18] Chuvil’deev VN, Petryaev AV. Phys Metals Metallogr 2000;89:122.